1

Chapter 1 Introduction

1-1 A brief history of Gallium Nitride epitaxial layers

Electroluminescence is a phenomenon in which a material emits light when electricity is

passed through it. Electroluminescence was one of the greatest discoveries of the twentieth

century and is rivaled by other forms of luminescence, including incandescence,

chemiluminescence, cathodoluminescence, triboluminescence, and photoluminescence.

Electroluminescence generally involves the production of light by a material without

producing heat. Electroluminescence differs from black body light emission, which is only

produced by mechanical actions. Electroluminescence is produced when an electric current is

passed through a semiconductor with tiny holes. As excited electrons pass over these holes,

they are emitted into the free air as particles known as photons. The semiconductor is usually

made of a film, whether organic or inorganic, and can be dyed by some types of metal, known

as a dopant. This is the most basic principle of all LEDs and most electronic lights in general.

Specifically, there is a phosphorous-based electroluminescent backlight in virtually all

electronic devices that is used to illuminate the Liquid Crystal Display (LCD).

Electronic lighting technology made its appearance with the advent of semiconductors

during the second part of twentieth century. The achievement of high-efficiency

eletroluminescence from semiconductor devices in early 1963 led to light-emitting-diode

(LED) technology [1-2]. Semiconductors with electroluminescence are expected a main part

in many optoelectronic systems, such as high power energy source system, optical memory

system, full color scanners or high definition ROM disk systems. Recently many researchers

have been done on electroluminescence, especially in Sinc selenide (ZnSe) [3], Silicon

carbide (SiC) [4], and Gallium Nitride (GaN) [2].

Electroluminescence has several advantages that has led it to be used so widely in

electronic devices. For example, electroluminescence requires very little current in order to

2

work. This is advantageous in small electronic devices because they often run on batteries that

are 1.5 Volts or smaller. Electroluminescent devices can be used to produce any color and are

very simple in design. Though electroluminescence can be advantageous in most cases, it

does have a few disadvantages. For example, electroluminescent devices require low amounts

of current but high amounts of voltage, usually between 60 and 600 Volts. In

electroluminescent devices that are connected to a power line, these voltages can be provided

consistently at any time of day. In battery-operated, handheld devices, however, the voltage

must be provided by a converter circuit that is built into the device.

1-2 The Introduction of Light-Emitting Diodes

Light-emitting diodes (LEDs) are a typical p-n junction device used under a forward bias.

The basic operating mechanisms are based on the electrical and optical properties of p-n

junction and of semiconductor materials. When the diode is forward biased, electrons are able

to recombine with holes and energy is released in the form of light. This effect is called

electroluminescence. As shown in Fig. 1.1[5], the wavelength of the light emitted, and

therefore its color, depends on the band gap energy of the materials forming the p-n junction.

Most commercially available LEDs are made from III-V compound semiconductors. Some

II-VI compound semiconductors such as ZnS and ZnSe are used in a few LEDs emitting

visible light, through these materials are not used frequently because of the difficulty of p-n

junction formation[6].

1-2-1 Historical Introduction

Electroluminescence was discovered in 1907 by the British experimenter H. J. Round of

Marconi Labs, using a crystal of silicon carbide and a cat’s-whisker detector[7]. Russian Oleg

Vladimirovich Losev independently created the first LED in the mid 1920s[8]; his research

was distributed in Russian, German and British scientific journals, but no practical use was

made of the discovery for several decades. LED development began with infrared and red

3

devices made with GaAs. Rubin Braunstein of the Radio Corporation of America first

reported on infrared emission from GaAs and other semiconductor alloys in 1955[9].

Braunstein observed infrared emission generated by simple diode structures using GaSb,

GaAs, InP, and SiGe alloys at room temperature and at 77 K. Experimenters at Texas

Instruments, Bob Biard and Gary Pittman, found in 1961 that GaAs gave off infrared light

when electric current was applied. Biard and Pittman were able to establish the priority of

their work and received the patent for the infrared light-emitting diode. Nick Holonyak Jr. of

the General Electric Company, who is seen as the “father of the light-emitting diode”,

developed the first practical visible-spectrum LED in 1962. Since then, related researches

went on continually.

In 1972, J. I. Pankove et al. fabricated the first blue LED using III-nitrides materials with

a Metal-i-n structure [10]. However, the device performance was limited by the poorly

conducting p-type GaN. The first high-brightness blue LED was demonstrated by Shuji

Nakamura of Nichia Corporation and was based on critical developments in GaN nucleation

on sapphire substrates and the demonstration of p-type doping of GaN which were developed

by Isamu Akasaki and H. Amano in Nagoya using a low-temperature buffer layer and

Low-Energy Electron Beam Interaction (LEEBI) technique, respectively[11]. In 1995,

Alberto Barbieri at the Cardiff University Laboratory investigated the efficiency and

reliability of high-brightness LEDs and demonstrated a very impressive result by using a

transparent contact made of indium tin oxide (ITO) on AlGaInP/GaAs LED. The existence of

blue LEDs and high efficiency LEDs quickly led to the development of the first white LED,

which employed a Y3Al5O12:Ce, or “YAG”, phosphor coating to mix yellow light with blue to

produce light that appears white. Nakamura was awarded the 2006 Millennium Technology

Prize for his invention.

4

1-2-2 Advantages of the LED Lighting

LEDs offer numerous benefits due to their mode of operation:

(1) Energy Efficiency

LEDs are highly efficient. In traffic signal lights, a strong market for LEDs, a red

traffic signal head that contains 196 LEDs draws 10W versus its incandescent

counterpart that draws 150W. Various estimates of potential energy savings range

from 82% to 93%. With the red signal operating about 50% of the day, the complete

traffic signal unit is estimated to save 35-40%. An estimation had been made that

replacing incandescent lamps in all of America’s some 260,000 traffic signals (red,

green and yellow) could reduce energy consumption by nearly 2.5 billion kWh. At

the end of 1997, more than 150,000 signals were retrofitted.

(2) Long Life

Some LEDs are projected to produce a long service life of about 100,000 hours. For

this reason LEDs are ideal for hard-to-reach/maintain fixtures such as exit sign

lighting and, combined with its durability, pathway lighting. This service life can be

affected by the application and environmental factors, including heat and if being

overdriven by the power supply.

(3) Range of Colors

LEDs are available in a range of colors, including white light. Fig. 1.2 shows the

chromaticity diagram[12]. White light can be produced through color mixing, like

red/blue/green, blue/yellow, or green/yellow-green/orange/purple. In addition,

through the innovative combination of various-colored LEDs, dramatic

color-changing effects can be produced from a single fixture through dynamic

activation of various sets of LEDs. Manufacturers such as Color Kinetics offer

fixtures that employ this principle. They offer track, theatrical, underwater, outdoor

and other fixtures utilizing variable-intensity LEDs that can provide more than 16.7

5

million colors, including white light. These fixtures can be individually controlled via

a PC, DMX controller or proprietary controller to generate effects including fixed

color, color washing, cross fading, and random color changing.

(4) Cool Light

In contrast to most light sources, LEDs radiate very little heat in the form of IR that

can cause damage to sensitive objects or fabrics. Wasted energy is dispersed as heat

through the base of the LED.

(5) Durable

LEDs are highly rugged. They feature no filament that can be damaged due to shock

and vibrations. They are subject to heat, however, and being overdriven by the power

supply.

(6) Small Size/Design Flexibility

A single LED is very small and produces little light overall. However, this weakness

is actually its strength. LEDs can be combined in any shape to produce desired lumen

packages as the design goals and economics permit. In addition, LEDs can be

considered miniature light fixtures; distribution of light can be controlled by the

LEDs’ epoxy lens, simplifying the construction of architectural fixtures designed to

utilize LEDs. A controller can be connected to an LED fixture to selectively dim

individual LEDs, resulting in the dynamic control of distribution, light output and

color. Finally, DC power enables the unit to be easily adaptable to different power

supplies.

(7) Dimming

LEDs can very easily be dimmed either by Pulse-width modulation or lowering the

forward current.

(8) On/Off Time

LEDs light up very quickly. A typical red indicator LED will achieve full brightness

6

in microseconds. LEDs used in communications devices can have even faster

response times.

(9) Cycling

LEDs are ideal for use in applications that are subject to frequent on-off cycling,

unlike fluorescent lamps that burn out more quickly when cycled frequently, or HID

lamps that require a long time before restarting.

1-2-3 Applications

Many application of LEDs are very diverse but fall into three major categories: Visual

signal application where the light goes more or less directly from the LED to the human eye,

to conveyca message or meaning. Illumination where LED light is reflected from object to

give visual response of these objects. Finally LEDs are also used to generate light for

measuring and interacting with processes that do not involve the human visual system. These

three categories are introduced below:

(1) Signals and Large Displays

The low energy consumption, low maintenance and small size of modern LEDs has

led to applications as status indicators and displays on a variety of equipment and

installations. Large area LED displays are used as stadium displays and as dynamic

decorative displays. Thin, lightweight message displays are used at airports and

railway stations, and as destination displays for trains, buses, trams, and ferries. The

single color light is well suited for traffic lights and signals, exit signs, emergency

vehicle lighting, ships’ lanterns and LED-based Christmas lights. Red or yellow

LEDs are used in indicator and alphanumeric displays in environments where night

vision must be retained: aircraft cockpits, submarine and ship bridges, astronomy

observatories, and in the field, e.g. night time animal watching and military field use.

Because of their long life and fast switching times, LEDs have been used for

automotive high-mounted brake lights and truck and bus brake lights and turn signals

7

for some time. The use of LEDs also has styling advantages because LEDs are

capable of forming much thinner lights than incandescent lamps. The significant

improvement in the time taken to light up (perhaps 0.5s faster than an incandescent

bulb) improves safety by giving drivers more time to react. Due to the relative

cheapness of low output LEDs, they are also used in many temporary applications

such as glowsticks and throwies and Lumalive, a photonic textile, artist have also

used LEDs for LED art.

(2) Lighting

With the development of high efficiency and high power LEDs, it has become

possible to incorporate LEDs in lighting and illumination. Replacement light bulbs

have been made as well as dedicated fixtures and LED lamps. LEDs are used as

street lights and in other architectural lighting where color changing is used. The

mechanical robustness and long lifetime is used in automotive lighting on cars,

motorcycles and on bicycle lights. The lack of IR/heat radiation makes LEDs ideal

for stage lights using banks of RGB LEDs that can easily change color and decrease

heating from traditional stage lighting, as well as medical lighting where IR-radiation

can be harmful. Since LEDs are small, durable and require little power, they are used

in hand held devices such as flashlights. LED strobe lights or camera flashes operate

at a safe, low voltage, as opposed to the 250 volts commonly found in xenon

flashlamp-based lighting. This is particularly applicable to cameras on mobile phones,

where space is at a premium and bulky voltage-increasing circuitry is undesirable.

LEDs are used for infrared illumination in night vision applications including

security cameras. A ring of LEDs around a video camera, aimed forward into a

retroreflective background, allows chroma keying in video productions. LEDs are

used for decorative lighting as well. Uses include but are not limited to

indoor/outdoor decor, limousines, cargo trailers, conversion vans, cruise ships, RVs,

8

boats, automobiles, and utility trucks. Decorative LED lighting can also come in the

form of lighted company signage and step and aisle lighting in theaters and

auditoriums

(3) Non-Visual Applications

Light has many other uses besides for seeing. LEDs are used for some of these

applications. The uses fall in three groups: Communication, sensors and light matter

interaction. The light from LEDs can be modulated very fast so they are extensively

used in optical fiber and free space optics (FSO) communications. This includes

remote controls, such as for TVs and VCRs, where infrared LEDs are often used.

Many sensor systems rely on light as the main medium. LEDs are often ideal as a

light source due to the requirements of the sensors. LEDs are used as movement

sensors, for example in optical computer mice. Touch sensing: Since LEDs can also

be used as photodiodes, they can be used for both photo emission and detection. This

could be used in for example a touch-sensing screen that register reflected light from

a finger or stylus. Many materials and biological systems are sensitive to, or

dependent on light. Grow lights use LEDs to increase photosynthesis in plants and

bacteria and vira can be removed from water and other substances using UV LEDs

for sterilization. Other uses are as UV curing devices for some ink and coating

applications as well as LED printers.

1-3 III-V Nitride Based LEDs

Creating sources of white light is the ultimate goal of solid-state lighting technology. The

most challenging application for LEDs is the replacement of conventional incandescent

andprobably, even fluorescent lamps. Attempts to up-convert long-wavelength emission of IR

LEDs to broader visible spectra were undertaken years ago by Galginaitis and Fenner (1968)

[13-14]. Berggren et al. (1994) demonstrated white light-emitting devices made from

electroluminescent organic semiconductors[15]. However, practical white LEDs became

9

feasible only after the development of high brightness blue AlInGaN emitters (Nakamura and

Fasol 1997) [16]. Based on short-wavelength LEDs, white LEDs that exploit the mixture of

two or three colors (dichromatic and trichromatic LEDs, respectively) are being developed.

The bandgap energy of AlGaInN varies between 6.2 and 1.95 eV depending on its

composition at room temperature, as shown in Fig. 1.1. Therefore, these III-V nitride

semiconductors are useful for light emitting devices especially in the short wavelength region.

Among AlGaInN system, GaN has been most intensively studied. GaN has bandgap energy of

3.4 eV at room temperature. Previous research on III-V nitrides has paved the way for the

realization of high quality crystals of GaN, AlGaN and InGaN, and of p-type conduction in

GaN and AlGaN[17-18]. The beginning of the growth of good quality epilayers was made by

Yoshida et al. [19]. They showed in 1983 that if an AlN buffer layer is grown by two step

method between the GaN film and the sapphire substrate, the quality of the layers improves.

Amano et al. [20] were also the first to obtain p-type conductivity in GaN which used Mg as

an accepter. In 1992 van Vechten et al. [21] suggested that Mg-H complexes are formed in the

as grown Mg-doped GaN and therefore hydrogen passivates the Mg acceptors. Later it was

found that annealing the Mg-doped GaN layers at >750 °C in nitrogen or vacuum also

activates the Mg acceptors[22]. By middle 1990s so much knowledge and experience in

technology had accumulated that the progress in designing and fabricating devices became

very rapid. High brightness blue LEDs have been fabricated on the basis of these results, and

luminous intensities over 1 cd had been achieved in 1994[23-25].

1-4 The development of Gallium Nitride epitaxial layers

The development of GaN-based devices has challenged the conventional thinking on the

material requirements for successful device fabrication in several ways. With the exceptions

of silicon on sapphire and GaN, no semiconductor has been commercialized exclusively using

heteroepitaxial materials; all other semiconductors have employed bulk substrates. Before

pn-junction GaN-based LEDs on sapphire substrates were demonstrated, good luminescence

10

efficiency, regardless of the semiconductor, was thought to require very low dislocation

densities, less than 106 cm-2. Researchers were surprised to measure excellent luminescence

efficiency from GaN LEDs despite dislocation densities four orders of magnitude higher.

Unlike the majority of semiconductors, in GaN dislocations did not seem to degrade its

optical and electrical properties. That sapphire was, and remains, the most common choice for

GaN based LEDs is particularly surprising, given its properties are seemingly unsuitable

based on the usual assumptions made in choosing a substrate for epitaxy; it has large lattice

constant and thermal expansion coefficient mismatches with GaN. [5]

Interest has always been high in the production and characterization of wide-bandgap

semiconductor materials for a wide range of short-wavelength and hightemperature

applications. Candidate material systems include silicon carbide (SiC), the Group II-VI

semiconductors zinc sulfide (ZnS) and zinc selenide (ZnSe), and the Group IIIA nitrides. The

hexagonal polytypes of SiC, although once considered excellent candidates for blue LEDs and

more recently for high-temperature and high-power electronic devices, are not suitable for

ultraviolet (UV) detector or emitter applications because of their indirect and comparatively

low bandgaps (Eg < 3.2 eV).

Recent progress[26] in developing blue–green LEDs and laser diodes in the ZnSxSe1–x

alloy system1 has been restricted to high-selenium alloys with bandgaps less than 2.8 eV.

Higher-bandgap alloys continue to be plagued by autocompensation effects, whereby the

introduction of a dopant atom is compensated by native defects and precludes well-behaved

conductivity control. Thus, the Group III-V nitrides are the preferred candidates for

short-wavelength applications. InN, GaN), and AlN have direct bandgaps of 2.0, 3.4, and 6.2

eV, respectively, with corresponding cutoff wavelengths of 620, 365, and 200 nm. Since they

are miscible with each other, these nitrides form complete series of indium gallium nitride

(In1–xGaxN) and aluminum gallium nitride (AlxGa1–xN) alloys, and it should be possible to

develop detectors with wavelength cutoffs anywhere in this range.[26]

11

1-5 An Overview of the Dissertation

The primary contents of this dissertation are subjects which are relating to the

mechanical properties and nanotribological characterization of GaN epitaxial layers on

different axis sapphire substrate, including the contact-induced deformation behaviors, the

repetition pressure-induced impairment behaviors, and nanotribological characterization

behaviors.

There five chapters in this dissertation. In chapter 1, a brief history of the progress of

lighting devices is presented. We also make a complete introduction about the LED, which

contains history, advantages, applications, development of III-V nitride-based LEDs, and

several key issues we might experience while studying it.

Chapter 2 reviews the GaN films including its characteristics, growth methods, and

applications. Besides, the nanoindentation and nanoscratch principle and technique were

introduced.

Chapter 3 introduces the experimental procedure and apparatus. The designed deposition

parameters for different axis structures of GaN films were exhibited. Finally, the detail

parameters of nanoindentation and nanoscratch for measuring mechanical properties and

nanotribological characterization were shown, respectively.

Chapter 4 analyzes mechanical properties and nanotribological characterization of GaN

films qualitatively and quantitatively. The ‘pop-in’ event was explained by the interaction of

the deformed region, produced by the indenter tip, with the inner threading dislocations in the

GaN films. In addition, the nanotribological behavior and deformation characteristics of the

GaN films were investigated by nanoscratch technique and atomic force microscopy (AFM).

Finally, chapter 5 summarizes the major results of this study and points out several

research areas which need further attention.

Figur

their

re 1.1 Band

lattice cons

dgap energy

stant[5].

y of various

12

materials fo

or visible emmission devvices as a fuunction of

Figur

red/b

re 1.2 Chro

blue/green, b

omaticity d

blue/yellow

diagram. Wh

w, or green/y

13

hite light c

yellow-gree

can be prod

en/orange/pu

duced throu

urple[12].

ugh color mmixing, likee

14

Chapter 2 Literature review

2-1 Introduction

Gallium nitride is a highly attractive material of groups III–V nitride semiconductors,

because it is great potential for the development of optoelectronic devices in blue/green light

emitting diodes, semiconductor lasers, and optical detectors. The fabrication of optoelectronic

devices based on GaN films requires to be understood in the mechanical properties as well as

its optical and electrical properties [22, 27–30], that is to say, the mechanical damages of GaN

films, such as film cracking and interface delamination caused by thermal stresses, chemical–

mechanical polishing, usually suppress the processing yield and application reliability of

microelectronic devices.

The heteroepitaxy using typical substrates (e.g., sapphire) with high lattice mismatch is

discussed in early report. GaN films on sapphire substrates exhibits large lattice mismatch

(about 14.5%) causing in-plane tensile strain in the GaN layers. In fact, the most common

orientation of sapphire used for GaN is the c-axis sapphire. And, the lattice mismatch of the

GaN films on a-axis [11 2 0] sapphire is less (2%) than that on c-axis (0 0 0 1) sapphire

(13.9%), which promises for excellent quality of GaN growth with improved surface

morphology [31]. Here, compared to bulk single crystals, the deformation properties of thin

films can serve strongly correlate to the geometrical dimensions and the materials defect

structure. The misfit dislocations at the interface play an important role such as carrier

mobility, and luminescence efficiency.

15

2-2 Physical properties, and development of GaN films

2-2-1 Physical properties of GaN films

The physical properties of GaN make it an attractive semiconductor for many electronic

[32] and optoelectronic devices [33]. Its wide, direct energy band gap makes it suitable for

short wavelength emitters (LEDs and diodes lasers) and detectors. The wide energy band gap

and good thermal stability of GaN is also advantageous for high-temperature and high power

electronics. Gallium nitride forms solid solutions with AlN and InN, making a wide range

(1.9–6.2 eV) of energy band gaps possible. This ability to form alloys is essential for

producing specific wavelengths for emitters, and for creating heterojunctions with potential

barriers into the device structures. Heat dissipation in devices is facilitated by GaN’s high

thermal conductivity compared to silicon and gallium arsenide. Both n- and p-type

conductivities are possible in GaN. Because the group III nitrides have a noncentrosymmetric

structure and significantly ionic chemical bonding, they are strongly piezoelectric, and also

undergo spontaneous polarization [34-35]. These effects can be advantageously employed, to

increase the sheet carrier concentration in heterostructure transistor [36]. A summary of the

physical properties of gallium nitride is given in Table 1. The thermal expansion coefficients

of semiconductors including GaN depend on the temperature, and here only changes in lattice

constants for certain temperature range are listed in this paper. The temperature-dependent

data of thermal expansion coefficient for various substrates including GaN, AlN, 6H-SiC,

3C-SiC, GaAs, Al2O3, ZnO, and MgO can be found in [42].

Gallium nitride normally has a wurtzite structure, with the space group of P63mc (no.

186). The wurtzite structure consists of alternating biatomic close-packed (0 0 0 1) planes of

Ga and N pairs stacked in an ABABAB sequence. Atoms in the first and third layers are

directly aligned with each other. Fig. 2.1 displays the perspective views of wurtzite GaN

along [0 0 0 1], [11 2 0] and [1 0 1 0] directions, where the large circles represent gallium

atoms and the small circles nitrogen. The close-packed planes are the (0 0 0 1) planes. The

16

group III nitrides lack an inversion plane perpendicular to the c-axis, thus, crystals surfaces

have either a group III element (Al, Ga, or In) polarity (designated (0 0 0 1) or (0 0 0 1)A) or

a N-polarity (designated (0 0 0 1) or (0 0 0 1)B). An excellent review on crystal polarity is

given by Hellman [48].

The zincblende structure (see Table 2.2 the space group is F43m) of GaN can be

stabilized in epitaxial films. The stacking sequence for the (1 1 1) close-packed planes in this

structure is ABCABC. Perspective views of the zincblende structure are shown in Fig. 2.2

Many different techniques are employed to assess the quality of GaN heteroepitaxial films. A

summary of some of the most important techniques with average, typical values for a thin (

17

2-2-2 The development of Gallium Nitride films

The development of GaN-based devices has challenged the conventional thinking on the

material requirements for successful device fabrication in several ways. With the exceptions

of silicon on sapphire and GaN, no semiconductor has been commercialized exclusively using

heteroepitaxial materials; all other semiconductors have employed bulk substrates. Before

pn-junction GaN-based LEDs on sapphire substrates were demonstrated, good luminescence

efficiency, regardless of the semiconductor, was thought to require very low dislocation

densities, less than 106 cm-2 [49-50]. Researchers were surprised to measure excellent

luminescence efficiency from GaN LEDs despite dislocation densities four orders of

magnitude higher. Unlike the majority of semiconductors, in GaN dislocations did not seem to

degrade its optical and electrical properties. That sapphire was, and remains, the most

common choice for GaN based LEDs is particularly surprising, given its properties are

seemingly unsuitable based on the usual assumptions made in choosing a substrate for epitaxy;

it has large lattice constant and thermal expansion coefficient mismatches with GaN.

Interest has always been high in the production and characterization of wide-bandgap

semiconductor materials for a wide range of short-wavelength and hightemperature

applications. Candidate material systems include silicon carbide (SiC), the Group II-VI

semiconductors zinc sulfide (ZnS) and zinc selenide (ZnSe), and the Group IIIA nitrides. The

hexagonal polytypes of SiC, although once considered excellent candidates for blue LEDs and

more recently for high-temperature and high-power electronic devices, are not suitable for

ultraviolet detector or emitter applications because of their indirect and comparatively low

bandgaps (Eg < 3.2 eV).

Recent progress in developing blue–green LEDs and laser diodes in the ZnSxSe1–x alloy

system1 has been restricted to high-selenium alloys with bandgaps less than 2.8 eV.

Higher-bandgap alloys continue to be plagued by autocompensation effects, whereby the

introduction of a dopant atom is compensated by native defects and precludes well-behaved

18

conductivity control. Thus, the Group IIIA nitrides are the preferred candidates for

short-wavelength applications. Indium nitride (InN), gallium nitride (GaN), and aluminum

nitride (AlN) have direct bandgaps of 2.0, 3.4, and 6.2 eV, respectively, with corresponding

cutoff wavelengths of 620, 365, and 200 nm. Since they are miscible with each other, these

nitrides form complete series of indium gallium nitride (In1–xGaxN) and aluminum gallium

nitride (AlxGa1–xN) alloys, and it should be possible to develop detectors with wavelength

cutoffs anywhere in this range.[26, 49, 51]

Improvements in the growth and characteristics of gallium nitride[52] have resulted in

significant inroads in device technology.[53] For example, high-efficiency blueemitting metal

insulator semiconductor LEDs[54] have been fabricated, and, with the advent of

magnesiumdoped p-type material, violet-emitting pn-junction LEDs[55] have been developed.

Prototype GaN metal semiconductor field-effect transistors have also been reported.[56] In

addition, high-quality, low-aluminumcontent GaN/AlxGa1–xN heterostructures have been

produced that exhibit quantum confinement effects by photoluminescence,[57] stimulated

emission under photon pumping,[58] and electron mobility enhancement.[59] Violet-emitting

In1–xGaxN/GaN LEDs with external quantum efficiencies of 0.22% have been made,[60] as

have double-heterostructure blue- and green-emitting In1–xGaxN/AlyGa1–yN LEDs with

external quantum efficiencies as high as 2.7%.[61] High-brightness blue, green, and yellow

LEDs based on AlxGa1–xN/In1–xGaxN quantum well structures have recently been reported.[62]

Finally, stimulated emission by current injection from an AlxGa1–xN/GaN/In1–xGaxN device

has been observed.[63-64]

19

Along with this progress in light-emitting and microwave devices, a small but significant

interest has grown in the development of photodetectors using the AlxGa1–xN alloy system.

Khan et al.[65] developed a metal–semiconductor–metal detector based on insulating GaN

and obtained a responsivity of 2000 A/W (amperes of photocurrent per unit incident photon

power) at 365 nm under an applied bias of 5 V. They recently reported an improved

responsivity of 3000 A/W by employing an enhanced mobility GaN/AlxGa1–xN

heterostructure.[66] However, their detector performance varied from sample to sample,

depending on the quality of the material. One main problem associated with GaN-based

detectors has been their slow response times, originating from the high defect density of the

starting material, which can also directly affect the stability and linearity of the devices.

A semiconductor-based photoconductive detector can be viewed as a radiation-sensitive

resistor,[67] as shown schematically in Fig. 2.3. In such a resistor, the current is generated by

incoming photons that excite electrons from the valence to the conduction band and by at least

some carriers being collected by one of the electrodes. Simply put, the photon-induced current

Iph is given by

gAq phI , (1)

where

q = electronic charge,

= quantum efficiency (i.e., number of carriers generated per incident photon in the primary

absorption process),

A = device area exposed to incident radiation,

= incident flux, and

g = photoconductive gain.

20

The two parameters influenced by the material characteristics are the quantum efficiency

and the photoconductive gain.

In its simplest form, the photoconductive gain is given by

ttt

g , (2)

where t is the majority carrier’s lifetime, and tt is the transit time of the carrier between

electrodes. This transit time can be approximated by

beVm

ltt , (3)

21

2-3 Formation of Gallium Nitride films

Because bulk gallium nitride crystals are not commercially available, most researchers

have relied on heteroepitaxy, which is crystal growth on substrates of another material, for

device fabrication. Most often, the lattice constant mismatch has been the primary criteria for

determining the suitability of a material as a substrate for GaN epitaxy. A wide variety of

materials have been studied for GaN epitaxy including insulating metal oxides, metals, metal

nitrides, and other semiconductors. A summary of crystal structure and lattice constants of

some of the candidate substrate materials for GaN epitaxy is given in Table 3.

In practice, properties other than the lattice constants including the material’s crystal

structure, surface finish, composition, reactivity, chemical, thermal, and electrical properties,

are also important in determining its suitability as a substrate, as these greatly influence the

resulting properties of the epitaxial layer, sometimes in unexpected ways. The substrate

employed determines the crystal orientation, polarity, polytype, the surface morphology, strain,

and the defect concentration of the GaN film. Thus, the substrate properties may ultimately

determine whether the device achieves its optimal performance. The influence of the substrate

on the polarity and polarization of the group III nitride epitaxial layer is particularly important.

The chemical reactivity and the conditions required for good quality epitaxy depend on the

polarity of the crystal. In many cases, the substrate controls the crystal polarity and the

magnitude and sign (tensile or compressive) of the strain incorporated into the epitaxial layers,

and thereby the extent of the polarization effect. Considerable variations are possible using a

variety of epitaxial growth techniques, as evidenced from the case of sapphire, where epitaxial

GaN films of either polarity can be controllably produced (see the subsection on sapphire).

Nevertheless, the choice of substrate does provide limits on what can be done in subsequent

processing.

Thus far, the vast majority of substrates studied produce [0 0 0 1] oriented GaN, as this

orientation is generally the most favorable for growing smooth films. However, interest in

22

GaN epitaxial layers with other orientations is increasing to eliminate the polarization effects.

Such effects can be deleterious for some optoelectronic devices, causing red shifts in emission.

In addition, piezoelectric effects in quantum wells can cause a spatial separation of electrons

and holes, thereby decreasing the recombination efficiency [68].

2-3-1 Structure and properties of Sapphire substrate

Sapphire, single crystal aluminum oxide, was the original substrate used in Maruskas and

Tietjen’s pioneering study of GaN epitaxy by hydride vapor phase epitaxy (HVPE) in 1969

[73], and it remains the most commonly employed substrate for GaN epitaxy. The large lattice

constant(15%) mismatch of sapphire (Al2O3) with GaN leads to high dislocation density (1010

cm-2) in the GaN epitaxial film [73]. These high defect densities reduce the charge carrier

mobility, reduce the minority carrier lifetime, and decrease the thermal conductivity, all of

which degrade device performance. Sapphire’s coefficient of thermal expansion is greater

than GaN, thus, producing biaxial compressive stress in the layer as it cooled from the

deposition temperature. For thick films, the stress can cause both the film and the substrate to

crack [74]. The thermal conductivity of sapphire is low (about 0.25 W/cm K at 100 8C), thus,

it is relatively poor at dissipating heat compared to other substrate materials. The cleavage

planes of epitaxial GaN layer do not parallel to those of sapphire, making laser facet

formation difficult. Sapphire is electrically insulating, thus, all electrical contacts must be

made to the front side of the device, reducing the area available for devices and complicating

device fabrication. In addition, there is evidence that oxygen from the sapphire causes

unintentional doping in the GaN layer, raising its background electron concentration [75].

Sapphire has the space group of R3c (no. 167) and is mainly composed of ionic bonds. The

single crystal can be described by both rhombohedral unit cells with volume 84.929 Ǻ3 and

hexagonal unit cell with volume 254.792 Ǻ3 [76], which is displayed in Fig. 3. In the

rhombohedral unit cell, there are 4Al3+ ions and 6O2- ions, and 10 ions in total. In the

23

hexagonal unit cell, there are 12Al3+ ions and 18O2- ions, and 30 ions in total. The unit cell

described in terms of hexagonal Miller-Bravais indices consists of six close packed (0 0 0 1)

planes of O2- ions, sandwiching 12 planes of Al3+ ions that occupy two thirds of the available

octahedral voids created by the O2- ions. The unreconstructed basal c-plane perspective views

for both unit cells are shown in Fig. 2.4, where the polyhedra are the cell boxes. The faceting

of sapphire crystal is shown in Fig. 2.5. All common surfaces employed for GaN epitaxy

including the (0 0 0 1) and (1 1 0 0) are non-polar. Thus, the polarityof the GaN epitaxial film

is primarily controlled by the process conditions to which it is subjected. Table 5 shows the

physical, chemical, thermal, mechanical and optical properties of sapphire.

2-3-2 Growth mechanisms

Epitaxial films have been grown on substrates of several orientations, and the resulting

orientations are presented in Table 6. By far, the most common orientation of sapphire used

for GaN is the c-plane sapphire. GaN epitaxy on c-plane sapphire results in c-plane oriented

films, but with a 308 rotation of the in-plane GaN crystal directions with respect to the same

directions in the sapphire. The 308 rotation of the (0 0 0 1) nitride plane with respect to the

sapphire (0 0 0 1) occurs to reduce the lattice constant mismatch; the mismatch would be 30%

without this rotation. With a proper nitridation and an optimized LT buffer layer of either AlN

or GaN, very smooth GaN films can be obtained. The c-plane sapphire is very difficult to

cleave, especially with a smooth cleave plane surface on the edge of the GaN layer. This is

made worse by the 308 rotation of the AlN and GaN c-plane relative to the sapphire c-plane,

such that the natural cleavage planes of the nitride and sapphire are not aligned. Stocker et al.

[78] reported that c-plane sapphire cleaves along both the a- and m-planes, with roughly

parallel cleaving of the m- and a-planes in the GaN film respectively. Large steps (90 nm)

were produced on the surface of the m-plane of the GaN which were unsuitable for laser

diode fabrication, hence, cleaves along the a-plane of sapphire unsuitable. Much smaller

24

striations occurred on the a-plane of GaN, making it a more promising for fabricating lasers.

Cleaving c-plane sapphire along its m-plane to expose the a-plane in the GaN remains an area

of active research. GaN films grown on a-plane sapphire are also oriented in the [0 0 0 1]

direction, and have the advantage that they are easily cleaved, along r-plane. This is

potentially advantageous for forming edge-emitting lasers. Although the lattice mismatch for

GaN films on a-plane sapphire is less (2%) than on c-plane sapphire (13.9%), no significant

differences in GaN film quality have been reported between these two substrates. Nakamura

et al. [79] has reported a pulsed MQW laser with cleaved facet mirrors fabricated in GaN on

a-plane sapphire, which might be a promising for production of nitride-based lasers on a large

scale.

The r-plane of sapphire is unusual, as it is one of only a few substrates capable of

producing GaN or AlN films with orientations other than the c-axis normal to the substrate.

For either GaN or AlN epitaxy on r-plane (110 2) sapphire, the a-axis [1 1 2 0] orientation is

produced. This is the preferred orientation for AlN-based surface acoustic wave devices, as it

produces the electromechanical coupling coefficient. For on-axis sapphire substrates with this

orientation, the AlN films are universally rough, forming a peak and valley morphology,

bounded by (0110 )and (0 11 2) facets. The roughness generally increases and becomes more

severe as the film thickness increases. By using r-plane sapphire substrates misoriented with a

48 tilt toward the [1101] negative a-direction, Shibata et al. [80] improved the quality of AlN

films by suppressing the formation of inverted twins.

From the viewpoints of lattice mismatch and crystal symmetry, (1 0 1 0) sapphire

(m-plane) seems the most suitable for GaN growth. However, the c-axis of GaN film grown

on a (1010)sapphire substrate has a non-zero inclination with respect to the c-axis of the

sapphire substrate. Thus, twins may be generated, and the structure and surface morphology

of GaN films becomes poor quality. This is a major disadvantage of a (1 0 1 0) plane

compared to a (0 0 0 1) plane.

25

Tripathy et al. [81] compared the quality of GaN layers with the same thickness (1.2 mm)

grown on c-, a-, r-, and m-plane sapphire substrates by MBE. GaN layers were smooth and

flat on cplane sapphire, somewhat rougher on a-plane, and rough on r-plane and m-plane

sapphire. Only (0 0 0 2) and (0 0 0 4) XRD diffraction peaks were seen from GaN on c-plane

and a-plane sapphire substrates due to the smooth surface morphology. The layers grown on

r-plane and m-plane had mixed orientations, with both (0 0 0 2) and (1 1 2 0) diffraction peaks.

From the smallest width of XRD rocking curve and narrow band PL spectrum, GaN on

c-plane sapphire suggested the best crystalline quality. Although the PL intensities are

comparable to that grown on c-plane sapphire, strong in-plane anisotropy of optical response

was found in GaN grown on a-, r-, and m-plane sapphire substrates. Matsuoka and Higiwara

[82] observed a significant improvement in the quality of GaN films deposited on m-plane

sapphire substrates misoriented by 15–208. This misorientation suppressed twinning and

enabled the growth of smooth, single crystal GaN films. Although the dislocation density was

still 50% higher than on c-plane sapphire, this is still a notable improvement over previous

efforts [82].

26

2-4 The nanoindentation technique for measuring mechanical properties of thin films

Indentation has been the most commonly used technique to measure the mechanical

properties of materials because of the ease and speed with which it can be carried out.

Traditional indentation testing involves optical imaging of the indent. This clearly imposes a

lower limit on the length scale of the indentation. During the past two decades, the scope of

indentation testing has been extended down to the nanometer range. This has been achieved

principally through the development of instruments capable of continuously measuring load

and displacement throughout an indentation [83-86]. A schematic representation to

nanoindenter apparatus is shown in Fig. 2.6.

Diamond is the most frequently used indenter material, because its high hardness and

elastic modulus minimize the contribution of the indenter itself to the measured displacement

[83]. For probing properties such as hardness and elastic modulus at the smallest possible

scales, the Berkovich triangular pyramidal indenter is preferred over the four-sided Vickers or

Knoop indenter because a three-sided pyramid is more easily ground to a sharp point [83, 86].

The shape is shown in Fig. 2.7. Another three-sided pyramidal indenter, the cube corner

indenter, can displace more than three times the volume of the Berkovich indenter at the same

load, thereby producing much higher stresses and strains in the vicinity of the contact and

reducing the cracking threshold. This makes this indenter ideal for the estimation of fracture

toughness at relatively small scales [87-88]. The spherical indenter initiates elastic contact and

then causes elastic–plastic contact at higher loads. This indenter, then is well suited for the

examination of yielding and work hardening.

The nanoindentations were sufficiently spaced (50 µm) to prevent from mutual

interactions. In order to obtain “film-only” properties, a commonly used rule of thumb

is to limit the indentation depth to less than 10 % of the films thickness [89-90]. The

deformation pattern of an elastic-plastic sample during and after indentation and a

27

typical load–displacement curve and are shown in Fig. 2.8 and Fig. 2.9, respectively.

Hardness and Young’s modulus were determined using the Oliver and Pharr analysis

[85]. Hardness (H) means the resistance to local plastic deformation of materials, which has

been conventionally obtained by measuring the projected contact area, Ac:

cAPH (4)

where P is the load. Elastic modulus E can be obtained from the contact stiff, using the

following relation:

cr AES 2 (5)

Ev

Ev

E ii

r

22 111

(6)

Where S is the experimentally stiffness, and β is a constant that depends on the geometry

of indenter and β=1.034 [91]. Er stands for the reduced modulus. Ei and vi are Young’s

modulus and Poission’s ratio of indenter, respectively. E and v are the same parameters for

the specimen. The load and stiffness are directly measured during an indentation,

contact area Ac and contact depth hc has relation 256.24 cc hA [92, 100]. By inserting

the calculated contacted area into Eqs.(5) and (6), the hardness and elastic modulus

were evaluated. For diamond which is the usual material of a Berkovich indenter, Ei = 1141

GPa and iv = 0.07 [89]. As the commonly done, we assume that v is 0.3.

28

2-5 The nanoscratch technique for measuring nanotribological characterization of thin films

Recently, improvements in the processing quality, quantity and variety of thin films have

led to increased interest in their potential as engineering materials and have increased the need

for a fullproperties. The latter is of interest since the generally high hardness values of thin

films would make them candidates for high wear applications. However, it has been argued

that hardness alone does not determine the wear resistance; in addition to resistance to

indentation, crack nucleation and propagation are also responsible for wear. Existing results of

wear studies often lead to conflicting reports of the wear resistance of thin films. This

discrepancy in wear data and its interpretation could be due to oversight of the importance of

small differences in processing conditions, which could generate differences in the quality and

structure of the alloys. One of the main reasons for the divergence in wear values is the wide

variety of acceptable wear tests, which can give significantly different results depending on

test conditions such as sliding vs. abrasive wear, sliding load and speed, and surface

roughness. The use of indentation to measure mechanical properties can be recalled to the

contact theory originally considered in the late 19th century by Hertz[93] and Boussinesq [94].

Hertz analyzed the problem of the elastic contact between two spherical surfaces with

different radii and elastic constants. His solutions form the basic theory in contact mechanics

and provide the analysis of non-rigid indenters. Boussinesq developed a method based on

potential theory for computing the stresses and displacements in an elastic body loaded by a

rigid axisymmetric indenter. His method has subsequently been used to derive solutions for

cylindrical and conical indenters. Another major contribution was made by Sneddon [95],

who derived general relationships among the load, displacement, and contact area for any

punch that can be described as a solid of revolution of a smooth function. Oliver and Pharr

[96] made a critical improvement to the indentation technique, who applied Snedden’s

formulation to determine the contact area at maximum load even if the contact area changes

29

during unloading. Their analytic procedure is to fit a power law function to the unloading

segment, which yields the contact stiffness as slope of this function at maximum load. This

slope is used to determinethe actual contact depth so that it is finally possible to derive the

indentation modulus and the indentation hardness.

In experiments, indentation tests were first performed by Brinell [83, 97-98], using

spherical and smooth balls from ball bearings as indenters to measure the plastic properties of

materials. The Brinell’s test was quickly adopted as an industrial test method soon after its

introduction and prompted the development of various macro and micro indentation tests [99].

During the past two decades, the scope of indentation tests have been extended down to the

nanometer range, such as the measurement of mechanical properties of nanocomposites and

nanometer thin films. This has been achieved principally through the development of

instruments capable of continuously measuring load and displacement throughout an

indentation. Oliver and Pharr [96] and Pharr [86, 100] proposed the significant technique to

determine the hardness and elastic modulus directly from the loading displacement curve

instead of measuring the indentation area. In recent years, Bhushan et al. [98, 101-102]

proved the mechanical properties are size-dependent in the nanoindentation, which has

received much attention. X. Li , Z.C. Li, and C. R. Taylor et al. [87-88, 103-105] used the

nanoindentation to estimate the fracture toughness of ultrathin films, which cannot be

measured by conventional indentation tests.

In addition, it is well known that the mechanochemical reaction of the contact surface

between the tip end and the surface plays an important role in nanomachining processes

because a friction force (or surface force) increases with the impact of a van derWaals force, a

meniscus force, and an electrostatic force on the nano/micro scale.Aborosilicate is stable

chemically because it mainly consists of silicate, whereas silicon easily combines with oxygen

or hydrogen. Accordingly, if a silicon–silicon bond is destructed by diamond tip sliding, oxide

or hydro-oxide is apt to be generated in reaction to oxygen or hydrogen in the atmosphere.

30

Specifically, Ando and Kaneko [106] revealed that a silicon protuberance formation is

deduced to be the mechanochemical reaction of silicon to diamond tip sliding through the

microfriction experiments [106-107]. They also have reported that the protuberance do not

occur under vacuum and dry nitrogen conditions. Miyake and Kim [107] have introduced

AFM nanoprotuberance processing technology for monocrystal silicon. They have controlled

the height of protuberances and the depth of grooves by varying the applied normal load and

diamond tip radius under ambient conditions, room temperature with humidity ranging

between 60 and 80%.[108]

Nanoscratch testing is a versatile tool for analysis of both thin films and bulk materials.

Nanoscratch provides the capability to investigate modes of deformation and fracture that are

not possible using standard indentation techniques. Nanoscratching is accomplished by

applying a normal load in a controlled fashion while measuring the force required to move the

tip laterally across the sample. By selecting the appropriate normal loading profile and lateral

displacement pattern, many different types of tests can be performed. The damage incurred

from the test is then typically observed using optical or SPM imaging. The schematic of the

nanoscratch machine is shown in Fig. 2.10, which basically includes a probe supported by a

2D transducer. The friction coefficient ( f ) is defined as

zx PPf (7)

where xP denotes the lateral load and zP indicates the normal load.[109]

Fig

Fig.

unit)

g.2.1 Persp

2.2 Perspec

; (b) [1 1 0]

pective view

ctive views

] (2 x2 x2 u

ws of wurtz2

of zincble

nits); (c) [1

31

ite GaN alo0]; (c) [1 0

ende GaN a

1 1] (2 x 2

ong various 1 0].[31]

along variou

x 2 units).

directions:

us direction

[31]

(a) [0 0 0 1

n: (a) [1 0 0

]; (b) [1 1

0] (1 x1 x1

PropEnerMaxi300 K77 KMaxiContn-typp-typMeltLattia (nmc (nm

Perce

ThermHeat ModuHardHardYield

erty gy band gap imum electroK

K imum hole mtrolled dopinpe pe ing point (Kce constants(

m) m)

entage chang

mal conductcapacity (30ulus of elasti

dness (nanoindness (Knoopd strength (1

Table 2.1

Fig. 2.3 A

(eV) (300 Kon mobility (

mobility (300ng range (cm-

) (300 K)

ge in lattice c

tivity (300 K00 K) (J/mol icity (GPa) ndentation, 3p, 300 K) (GP000 K) (MPa

Properties o

A diagram o

K) (cm2/V s)

0 K) (cm2/V s-3)

constants (30

) (W/cm K)K)

00 K) (GPa)Pa) a)

32

of GaN on

of a typical p

s)

00–1400 K)

)

sapphire (0

photocondu

Value 3.44 1350 19200 13 1016 to 4x1016 to 6x>2573 (a 0.3188430.518524

0/ aa

0/ cc2.1 35.3 210 ± 2315.5 ±0.910.8 100

0 0 1) subs

uctor circuit

x1020 x1018

at 60 kbar)

3 4

0.5749, 0.5032

9

strate.

.[67]

Reference[37]

[38]

[39]

[40]

[41]

[42]

[42]

[43] [44] [45] [46] [46] [47]

Tab.

Fig. 2rhom

Ma

Sem w z r w z r Z S S G G SOxi A

2.2 Substra

2.4 Perspecmbohedral u

aterial

miconductorsw-GaN zb-GaN r-GaN w-AlN zb-AlN r-AlN ZnO[69,70] -SiC SiC SiC GaAs GaP[71,72] Si ides and sulpAl2O3 (sapph

ate candidat

ctive views iunit cell; (b)

St

s WZiRoWZiRoW3C4H6HZiZiD

phides hire) Rh

es for GaN

in (2 2 1along the (

tructure

Wurtzite incblende ock salt

Wurtzite incblende ock salt

Wurtzite C (ZB) H (W) H (W) incblende incblende iamond

hombohedra

33

epitaxy.[31

1) unit cells(0 0 0 1) dir

Space gro

P63mc F43m Fm3m P63mc F43m Fm3m P63mc F43m P63mc F43m F43m Fd3m

al R3c

1]

s: (a) along tection in he

oup a 0.30.40.40.30.40.40.30.40.30.30.50.50.5 0.4

the [0 0 0 1exagonal un

Lattice cob

31885 4511 422 31106 438 404 32496 43596 3073 30806 56533 54309 54310

4765

] direction init cell. [31]

onstants (nm)c 0 0 0 11 1

in a

)

.5185

.49795

.52065

.0053

.51173

.2982

Fig. 2

Fig. load senso

2.5 Commo

2.6 A schemapplication

or [85].

on facets of

matic repren coil; (D)

sapphire cr

sentation otindentation

34

rystals: (a) v

t nanoindenn column

view down c

nter apparatguide sprin

c-axis; (b) s

tus. (A) samngs; (E) ca

surface plan

mple; (B) inapacitive di

nes. [77]

ndenter; (C)isplacement

) t

Fig. 2inden

2.8 The defntation. [85

F

formation p5,112]

Fig. 2.7 The

pattern of a

35

e standard B

an elastic-p

Berkovich tip

plastic sam

p. [113]

mple duringg and after

Fig. 2meas

2.9 Schemasured param

atic illustratmeters. [85,1

ion of inden112]

Fig. 2.10 S

36

ntation load

Sketch of n

d–displacem

anoscratch.

ment data sho

[109]

owing impoortant

37

Chapter 3 Experimental Apparatus and Procedures

3-1 Experimental procedures

In the experiment, the parameter and primary design are show in Fig. 3.1.

3-2 The process of the GaN films growth

3-2-1 Sapphire substrate preparation

The substrates used in experiments were 2-inch A-plane (1 1 2 0) and C-plane(0 0 0 1)

orientated sapphire wafer with thickness 420 ± 25um.

3-2-2 GaN films deposition

Sapphire wafer was introduced into load-lock chamber of metal-organic chemical vapor

deposition (MOCVD). To fabricate the GaN films, a 10-nm-thick AlN buffer layer was grown

on the sapphire substrate, and then both GaN films (thickness: ca. 2 um) were grown on top of

the buffer layer through MOCVD at 1100 oC, using triethylgallium (TEGa),

trimethylaluminum (TMAl), and ammonia (NH3) as the gallium, aluminum, and nitrogen

sources, respectively. The GaN films grown on c-plane (0001) and a-plane (11 2 0) sapphire

had a [0001] orientation and a [1 1 2 0] orientation.

3-3 Surface morphologies measurement

In this section, the measurement methods used to characterize the material of GaN grown

by ULVAC MBE system are described.

3-3-1 Field emission scanning electron microscope (FE-SEM)

The FE-SEM consists of a field emission electron source rather than a thermionic emission

source used in a Thermal Tungsten wire SEM. A FE-SEM is therefore has a cold cathode. This

SEM provides higher resolution imaging with higher beam density (brightness), and longer tip life.

The primary electrons enter a specimen surface with energy of 0.5 to 30 keV to generate many

low energy secondary electrons. The second electrons emission depends largely on the

accelerating voltage of the primary electrons and also the probing incident angle on the specimen

38

surface. In general, a large quantity of secondary electrons is generated from the protrusions and

the circumferences of objects on the specimen surface, causing them to appear brighter then

smooth portions. Under a constant accelerating voltage, an image of sample surface can thus be

constructed by measuring the secondary electron intensity as a function of the position of

scanning primary electron beam. In this study, JEOL-JSM7001F FE-SEM is used routinely to

check the GaN surface morphology as well as the film thickness.

3-3-2 Atomic force microscope (AFM)

AFM is used to check the surface topography of a material by measuring the interaction

force between the AFM‟s probe and surface structure. The AFM probe consists of a sensitive 19   

cantilever and a sharp tip (tip radius

39

penetration depth of 15, 40, and 300 nm. Each indentation was separated by 30 mm to avoid

possible interference between neighboring indents. Here, all indents were performed at room

temperature. The analytic method developed by Oliver and Pharr was adopted to determine

the hardness (H) and Young’s modulus (E) of the GaN films from the load–displacement

curves [85, 112]. In addition, the cyclic nanoindentation deformation behaviors of the GaN

films/a-axis sapphire substrate were conducted under force contact stiffness measurement

(FSM) technique with repetition pressure-induced impairment at the constant load of 50 mN.

In FSM measurement, each indentation location was separated by 30 μm in order to avoid

possible interference between neighboring indents. Hence, the CSM mode can obtain

hardness and Young’s modulus relative to the indentation depth and is used in this

experiment [85, 91, 113-114]. The instrument was calibrated by using a standard fused

silica sample before measuring the mechanical characterizations of GaN epitaxial layers.

The drift rate preset to < 0.05 nm/s before the beginning of each indentation.

Frequency of 45 Hz was used to avoid the sensitivity to thermal drift and loading

resolution was 50 uN [91, 115]. Ten indents were made on samples to minimize the

deviation of the results.

40

3-5 Nanotribological characterization measurement

The nanotribological properties of GaN films on A- and C-axis sapphire substrate were

determined by combining AFM (Digital Instruments Nanoscope III) together with a

nanoscratch measurement system (Hysitron), operated at a constant scan speed of 2 um s-1.

For the GaN epilayer/sapphire systems, constant forces of 2000, 4000, and 6000 μN were

applied. The maximum load was then maintained while forming 10-um-long scratches.

Surface profiles before and after scratching were obtained by scanning the tip at a 0.02-mN

normal load (i.e., a load sufficiently small that it produced no measurable displacement). After

scratching, the wear tracks were imaged using AFM.

3-6 Cathodoluminescence properties measurement

The cathodoluminescence properties of GaN epilayers with cyclic nanoindentation

deformation of the GaN epilayers/a-axis sapphire substrate were determined by

cathodoluminescence (CL) mapping. The cathodoluminescence (CL) mapping acquired at a

wavelength of 550 nm from an indented GaN films indented was characterized using a CL

apparatus (HORIBA Co., Ltd.). The room-temperature CL measurements were performed

using an electron beam energy of 20 keV. The CL light was dispersed by a 2400 nm grating

spectrometer and detected by a liquid N2-cooled charge-coupled device.

41

3-7 Experimental apparatus

(a) Metal-organic chemical vapor deposition (MOCVD)

MOCVD system features include:

1. Wafer size: 2".

2. Number of wafer: up to 6.

3. Stainless Steel or Quartz Chamber.

4. IR or RF heating.

5. Chamber temperature control: 1300

6. Wafer Rotation.

7. Pressure Control: 0 ~ 800 Torr.

8. Deposited materials: III-V. (GaN, GaAs, AlGaN, InP, etc)

(b) Field emission scanning electron microscope (FE-SEM)

JEOL-JSM7001F was used routinely to check the nanoindentation mark of GaN surface

morphology as well as the thin film.

(c) Atomic force microscope (AFM)

Digital Instruments DI 5000 was utilized to observe the 3D surface morphologies and

roughness of the GaN epilayers.

(d) Nano Indenter XP instrument

Nano Indenter XP instrument (MTS Cooperation, Nano Instruments Innovation Center,

TN, USA) was utilized to investigate the hardness and elastic modulus of the GaN epilayers.

(e) Hysitron Triboscope instrument

Hysitron Triboscope instrument was utilized to investigate the Nano Indent, Nano

Dynamic Mechanical Analyzer (Nano-DMA), Nano Scratch, and Wear test of thin film.

(f) Crygenic Cathodoluminescence system instrument

Crygenic Cathodoluminescence system instrument(JEOL, JSM7001F) to investigate the

cathodoluminescence properties of GaN films with cyclic nanoindentation deformation of the

GaN films/a-axis sapphire substrate.

42

Fig. 3.1 The parameter and primary design.

43

Chapter 4 Primary result

4-1 Investigation of contact-induced deformation behaviors of the GaN films/a-axis

sapphire substrate

4-1-1 Motivation

The mechanical damages of GaN films, such as film cracking and interface delamination

caused by thermal stresses, chemical–mechanical polishing, usually suppress the processing

yield and application reliability of microelectronic devices. The heteroepitaxy using typical

substrates with high lattice mismatch is discussed in early report. GaN films on sapphire

substrates exhibits large lattice mismatch (about 14.5%) causing in-plane tensile strain in the

GaN layers. In fact, the most common orientation of sapphire used for GaN is the c-axis

sapphire. And, the lattice mismatch of the GaN films on a-axis(11 2 0) sapphire is less (2%)

than that on c-axis (0 0 0 1) sapphire(13.9%), which promises for excellent quality of GaN

growth with improved surface morphology [31, 116]. Here, compared to bulk single crystals,

the deformation properties of thin films have strongly correlation to the geometrical

dimensions and the materials defectstructure. The misfit dislocations at the interface play an

important role such as carrier mobility, and luminescence efficiency. In this respect,

nanoindentation has proven to be a powerful technique for probing the information on the

mechanical properties of GaN films with characteristic dimensions in the sub-micron regime,

such as hardness and Young’s modulus [117-118]. During indentation load of GaN films, the

nanoindentation-induced discontinuity (so-called ‘pop-in’) in the load–displacement curve has

been observed in elsewhere investigation. The slip band movement [45, 119-120] and

dislocation nucleation mechanism [117,121] have been proposed to explain this ‘pop-in’ event.

Most of the studies were carried out on c-axis GaN films or bulk single crystals [85, 122-123].

Comparatively, the GaN films was observed in their mechanical damages, however the detail

work in the a-axis GaN films is still not yet unclear.

44

In this paragraph, significant investigation of the mechanical characterizations on GaN

films is motivated. The pressure-induced impairment of GaN films/a-axis sapphire substrate

has been investigated by using nanoindentation technique and atomic force microscope

(AFM).

4-1-2 Results and discussion

The load-indentation depth data for the grown GaN epilayrs by MOCVD with 2 µm

thickness on a-axis sapphire substrate is discussed, the multiple discontinuities during loading

were clearly revealed in Fig. 4.1(a). The hardness and elastic modulus as a function of

indentation depth can be obtained from the CSM measurements, as illustrated in Fig. 4.1(b)

and (c), respectively. The hardness and elastic modulus of the GaN films/a-axis sapphire

substrate are determined to be 15.9 ± 0.8 and 394.6 ± 12 GPa, respectively. Table 4.1

displays some of relevant data for comparison [119, 124, 129]. It is believed that the

discrepancies among the mechanical parameters obtained by various indentation methods are

mainly due to the specific tip–surface contact configuration and stress distribution inherent in

the different plane sapphire substrate. At beginning, the ‘pop-in’ event took place at a depth of

16 nm (at a critical load of 0.16 mN) was observed and, the contact stress has is 35.26 GPa

(see the inset in Fig. 4.1(b)). The critical contact stress of the penetration depth in 16 nm is

required to cause the initial ‘pop-in’ event in the a-axis GaN film. This case is relatively large

with that of GaN film on c-axis sapphire substrates (23 nm; at a load of 0.48 mN; contact

stress 22.5 GPa) [125]. The phenomenon suggested that the low indentation load (0.45 mN)

serve the Berkovich tip in the increasing stress on the film’s surface than the Berkovich tip on

the GaN film’s surface. In fact, the data obtained from Berkovich diamond indenter in our

result were also very large with that of spherical tip (140 nm; at a load of 30 mN; contact

stress 9 GPa) as reported in the other study [124]. Therefore, the initial sudden displacement

discontinuity generally accompany with maximum shear stress generated under the indenter

45

and, is more than that of the remaining other ‘pop-ins’. It is also observed that the

corresponding indentation depth versus hardness data from the hardness value increased

highly about 36 GPa before the ‘pop-in’ event at 16 nm (see the inset in Fig. 4.1(a)). The fact

was attributed from the sharper tip that induces plasticity at a shallow indentation depth.

Therefore, the hardness values can be calculated accurately [119]. It is suggested that the

‘pop-in’ events partially induce from the heteroepitaxial thin film, which can also be

attributed to the inner-defects of GaN films.

Evidently, due to the soft GaN film grown epitaxially on a hard sapphire substrate, it is

easy to find that the differences in lattice parameter and thermal expansion coefficients

between the film and substrate. This may introduce a misfit strain at the interface, and then

the strain is released by generating threading dislocations in the GaN films [126-127]. In the

experiment of indentation loading, the deformed region was punched out on the top edge

from the indenter tip; the threading dislocation may be to cause a sudden increase in plasticity,

therefore the sudden discontinuity in the load-indentation depth curve is showed.

Navamathavan et al. [125] concluded that the threading dislocation in the GaN films/c-axis

sapphire substrate can suddenly propagate after acquiring threshold energy from the deformed

region. This is the reason why the multi ‘pop-ins’ of the GaN films is occurred. However,

some of the threading dislocations terminate as they yet to reach the film surface [128]. At the

same time, the GaN films/a-axis sapphire substrate can be interacted by indenter tip to appear

deformed zone, the threading dislocations then excurse into films through certain depths,

afterward causing the ‘pop-in’ event. It is conjectured that the mechanism of the event is

caused by slip bands [119-120, 124], and/or dislocation nucleation [117].

In Fig. 4.2(a), the AFM image of the residual indentation mark is revealed that none of

crack and particle even occurred after the indentation depth. From the section view of the

indentation mark, Fig. 4.2(b) shows the residual volume from the edge of indentation.

Dislocation-induced ‘pop-in’ event tends to associated with two distinct deformation

46

behaviors before (pure elastic behavior) and after (elastoplastic behavior) the phenomenon. It

is speculated that the multiple ‘pop-in’ events are revealable over the indentation load and

penetration depth. It is suggested that the initial ‘pop-in’ event is owing to the role of

dislocation. Therefore, this could be explored that the occurrences of a discontinuity in the

load-indentation depth curve. For this case, the investigation of initial deformation is

measured under critical depth, which films deformed elastically and no residual deformation

is observed (Fig. 4.3(a)). At the same time, while the indentation is stopped after the exactly

‘pop-in’ distance, the residual mark depth is observed as shown in Fig. 4.3(b). This leads to

larger deviations in the penetration depth versus indentation load curves. We suggested that

the ‘pop-in’ event of the GaN films could be contributed from an entirely plastic deformation

case. The present results are similar with the work that the GaN films can be elastically

flexure before the ‘pop-in’ event [125].

From above results and discussion, the primary deformation mechanism of the GaN

films is due to dislocation nucleation and propagation along the easy slip systems. The

mechanisms of responsible for the dislocation recovery appears to be associated with the

activation of dislocations sources bring by loading-unloading cycle of the GaN films. The

plastic deformation prior to loading-unloading cycle is associated with the individual

movement of a small number of new nucleation, large shear stress is quickly accumulated

underneath the indenter tip. When the local stress underneath the tip reaches at high level

cycles, a burst of collective dislocation movement on the slip systems is activated, resulting in

a release of local stress [125]. The extensive interactions between the dislocations slipped

along the GaN surface, confined the slip bands and resulted in a ‘pop-in’ event due to the

deformed and strain-hardened lattice structure.

47

4-1-3 Conclusion

In this part of experimental results, the contact-induced deformation behaviors of the

GaN films/a-axis sapphire substrate were investigated by nanoindentation and AFM technique.

The deformation mechanisms of the GaN films result in a ‘pop-in’ event during

loading-unloading cycle, especially lead to deviations in the penetration depth versus

indentation load curves. The punching out of the threading dislocations beneath the deformed

region may contribute to the ‘pop-in’ event, which has been shown as one of the mechanisms

responsible for the plastic deformation of the GaN films.

48

4-2 Investigation of repetition pressure-induced impairment behaviors of the GaN

films on a-axis sapphire substrate

4-2-1 Motivation

In terms of GaN film, the mismatch of lattice constants and thermal expansion

coefficients in this heteroepitaxy serve high dislocation densities and high level of residual

strain in the postgrowth of thin film, which affects its mechanical properties and etc.[22,30,

61, 130-133]. The mechanical damages of GaN films, such as film cracking and interface

delamination caused by thermal stresses, usually suppress the processing yield and application

reliability of microelectronic devices. In this scenario, to access accounting precise in the

mechanical properties of GaN remains a challenge. It has been made to develop processes to

grow thick GaN films and subsequently separate them from their substrates as evidenced [31,

134-138]. Notice that, GaN film on sapphire substrates exhibits large lattice mismatch (about

13.9 %) causing in-plane tensile strain in the GaN film. The most common orientation of

sapphire used for GaN is the c-axis sapphire. In fact, the lattice mismatch of the GaN film on

a-axis (11 2 0) sapphire is less (2 %) than that on c-axis (0001) sapphire (13.9 %), which

promises for excellent quality of GaN growth with improved surface morphology [31].

Above of them, compared to bulk single crystals and the deformation mechanism of the

films can serve highly correlate to the geometrical dimensions and the materials

defect-structure. The misfit dislocations at the interface of GaN films, however, can play an

important role of carrier mobility and luminescence efficiency, etc. As a consequence,

nanoindentation has proven to be a powerful tools for probing the information on the

mechanical properties of GaN films with characteristic dimensions in the sub-micron regime,

such as hardness and Young’s modulus [117-118, 139-151]. In terms of indentation load on

GaN films, the nanoindentation-induced discontinuity in the load-displacement curve has

been revealed. The slip band movement [45, 119, 142-143] and dislocation nucleation

mechanism [117, 140, 144]to explain this pop-in event on c-axis GaN films or bulk single

49

crystals have been investigated [122, 145]. In our prior article, the proposal of a-axis GaN

films had been attempts in pop-in study [146], but for reasons that are unclear from repetition

pressure-induced impairment events.

In this paragraph, the pressure-induced impairment of GaN films/a-axis sapphire substrate

has been investigated by using nanoindentation, AFM, and CL mapping. A significant

investigation of the mechanical characterizations on GaN films is motivated here.

4-2-2 Results and discussion

In term of indentation, residual impression of repeated indent mark was examined by

tapping mode 3D-AFM to check for the evidence of slip, cracking, and pile-up/sink-in

phomenon. Shown in Fig. 4.4 are typical amplitude-mode AFM images of a (a) typical cycle,

(b) three, (c) six, and (d) nine loading/reloading cycles to maintain a maximum load of 25

mN. In Fig. 4.4(a), AFM image clearly illustrated that slightly activation of pyramidal slip

near the indent mark, which occurs at typical cycle nanoindentation, is suppressed in the GaN

films. A comparison of three and six loading/reloading cycles nanoindentation images [Fig.

4.4(b)-(c)] gives a gradation of protrusion that mean some pyramidal slip nucleation rather

than that of typical cycle one. It is suggested that the slip mechanism can be responsible for

the pile-up events during the loading of GaN. Furthermore, the AFM image of the residual

indentation mark at nine loading/reloading cycles is revealed that the pyramidal slip and

partially crack [Fig. 4.4(d)]. The dislocation-induced ‘pop-in’ event tends to be associated

with two distinct deformation behaviors before (pure elastic behavior) and after (elastoplastic

behavior) the phenomenon. The hardness (H) and elastic modulus (E) as a function of

indentation depth can be obtained from the FSM measurements; for the measured results of

the GaN films/A plane sapphire substrate, H were 16.4± 0.2, 15.6± 0.1, 16.6± 0.2, 26.5± 0.8

GPa, while E were 409.7± 12, 400.7± 11, 383.5± 9, 758.8± 18 GPa, respectively. The soft

GaN films was grown epitaxially on a hard sapphire substrate, it is tend to find that the

differences in lattice parameters and thermal expansion coefficients between the film and

50

substrate. It may introduce a misfit strain at the interface and the strain is released by

generating threading dislocations in the GaN films [126-127, 147-148]. From the plane view

of the indentation mark, Fig. 4.4 shows the residual volume from the edge of indentation.

There is massive material transfer from below the indented zone to the sides of the

indentation mark. Navamathavan et al. [149] revealed that the residual impression of the

indentation mark by Berkovich indenter tip. The pile-up event on the morphological surface

is clearly observation from SEM technique. Barsoum et al. also reported [150] that pop-ins

event and the cracks under a Berkovich indenter were observed. While the indenter suddenly

to penetrate into the basal planes, they must rupture [Fig. 4.4(d)]. Since the rupture is

associated with the repeated indenter cycles, the relatively large displacements at low stresses

below the bend are most probably related to the back and forth motion of delaminated. It is

believed that the discrepancies in the mechanical parameters obtained by repeated indentation

methods are mainly due to the specific tip–surface contact configuration and stress

distribution inherent in the system of GaN films/a-axis sapphire.

In Fig. 4.5, the CL mappings acquired at a wavelength of 550 nm from a GaN films

indented were presented at (a) typical, (b) three, (c) six, and (d) nine loading/reloading cycles,

respectively. These CL images reveal the distribution of indentation-induced extended defects;

the cube-corner indenter form and these extended defects reflect some of the radial symmetry

of the stress field [Fig. 4.5(a)-(c)]. In contrast, another luminescence/topography distribution

CL image revealed a crack line that formed after nine loading/reloading cycles [Fig. 4.5(d)].

Slip enhancement may contribute to the crack that can be observed during loading/reloading

and, the mechanism may be due to the plastic deformation [141]. In the GaN films, the

indented area below the propagation of dislocations is observed by CL image that also

revealed a detectable reduction in the intensity of the CL emission [148-151]. In the case of

six loading/reloading cycles, only a small percentage of indentations was plastically deformed

[Fig. 4.5(c)]. CL imaging could distinguish between the indentations that had undergone

51

several degree of plastic deformation, which location was purely in the elastic regime [Fig.

4.5(d)]. As a result, it is detected an observable CL impression only after the pile-up event,

providing convincing evidence that the phenomenon involves the nucleation of slip as the

deformation mode.

Above of them, the load-indentation behavious for the GaN films on the A plane

sapphire are revealed in Fig. 4.6. At beginning, the FSM serves the typical-load vs.

penetration depth of GaN film [Fig. 4.6(a)]. After the three cycles, the load vs. penetration

depth of GaN film is recorded; the repetition pressure-induced curve means gradually elastic

deformation. The repetition pressure-induced under three and six loading/reloading cycles

nanoindentation images [Fig. 4.6(b)-(c)] gives a more suffer in strain hardening compared

with that of typical one [Fig. 4.6(a)]. It is noted that the strain hardening is consist of the

decreased in the penetration depth. While Fig. 4.6(d) is superimposed, it becomes evident that

the crack and the increased penetration depth are almost identical at CL imaging [Fig. 4.5(d)].

From Barsoum et al. [118], noted that all slips are basal and the stress-strain curves both

before and after the pop-ins for both C and A orientations presumably were almost perfectly

superimposable in sapphire. It is thus reasonable to implicate basal slip for the deformation in

the A orientation.The corresponding penetration depth vs. repetition cycles concluded in Fig.

4.7; the strain hardening of GaN films partially induce as the penetration depth are observed

to decrease from typical to six loading/reloading cycles, but nine loading/reloading cycles.

Consequently, the system (GaN films/a-axis sapphire) can be interacted by indenter tip to

appear deformed zone, the threading dislocations then excursion into films through certain

depths, afterward causing the pop-in event occurred. It is conjectured that the mechanism of

the event is caused by slip bands [119, 124, 152-154], and/or dislocation nucleation [117,

155]. However, the inner failure of the GaN films is observed while crack is suddenly

occurred at the condition of nine loading/reloading cycles. It is speculated that the repetition

pressure-induced impairment during loading were contributed by the multiple pop-in events

52

and, are revealable over the indentation load and penetration depth [40, 44]. From the

previous reported [118, 156-157], they discussed that the deformation mechanisms in

freestanding films. The successive indentations in the same location generated a series of

hysteresis loops in which the first loop was slightly open, but subsequent loops were perfectly

superimposable. For this case, the investigation of initial deformation of GaN films is

measured under cyclic nanoindentation, which films deformed elastically and residual

deformation is observed. This leads to larger deviations in the penetration depth vs. cyclic

nanoindentation. The mechanisms of responsible for the dislocation recovery appears to be

associated with the activation of dislocations sources bring by loading-unloading cycle of the

nanoindentation